In recent years, production of Nd—Fe—B alloys serving as magnet alloys has sharply increased by virtue of excellent characteristics of the alloys, and these alloys are employed in HDs (hard disks), MRI (magnetic resonance imaging), a variety of motors, etc. Typically, Nd is partially substituted by another rare earth element such as Pr or Dy and Fe is partially substituted by another transition metal element such as Co or Ni. Such substituted alloys as well as Nd—Fe—B alloys are generally referred to as R—T—B alloys. Herein, R represents at least one rare earth element including Y, and T represents transition metals including Fe as an essential element. Fe may be partially substituted by Co or Ni. Other elements such as Cu, Al, Ti, V, Cr, Mn, Nb, Ta, Mo, W, Ca, Sn, Zr, and Hf may be added, to the R—T—B alloys, singly or in combination of two or more species. B represents boron and B may be partially substituted by C or N.
An R—T—B alloy contains, as the main phase, a ferromagnetic phase formed of R2T14B crystals, which contribute to magnetization, and a nonmagnetic R-rich phase having a low melting point and containing a rare earth element(s) at high concentration. Since the R—T—B alloy is an active metallic material, the alloy is generally melted and mold-cast in vacuum or under inert gas. In a typical method of producing a sintered magnet, an ingot of the alloy is pulverized to powder having a particle size of about 3 μm (as measured by means of FSSS (Fisher Sub-Sieve Sizer)); the powder is subjected to press-forming in a magnetic field; the resultant compact is sintered in a sintering furnace at a temperature as high as approximately 1,000 to 1,100° C.; and in accordance with needs, the sintered product is heated, mechanically processed, and plated for corrosion prevention.
The R-rich phase plays the following important roles in sintered magnets formed of the R—T—B alloy.
(1) Since the R-rich phase has a low melting point, the phase liquefies during sintering, thereby contributing to achievement of high density of the resultant magnet, leading to improved magnetization.
(2) The R-rich phase functions to smoothen grain boundaries, thereby reducing the number of nucleation sites of reversed magnetic domains, thereby enhancing the coercivity.
(3) The R-rich phase magnetically insulates the main phase, thereby enhancing the coercivity.
The distribution of the R-rich phase in a magnet, the final product, depends greatly on the microstructure of the as-cast alloy ingot. Specifically, when the alloy ingot is produced through mold casting, a slow cooling rate of the cast ingot often results in formation of large crystal grains in the R2T14B phase, and R-rich phase forms large aggregates which are locally present in the ingot. Thus, the particle size of the pulverized alloy ingot becomes considerably smaller than the grain size of crystals present in the R2T14B phase or than the size of dispersed R-rich phase. Therefore, particles formed only of the main phase (R2T14B phase) containing no R-rich phase and particles formed only of the R-rich phase are produced, whereby homogeneously mixing the main phase and R-rich phase becomes difficult.
Another problem involved in mold casting is that γ-Fe tends to be formed as primary crystals, due to the slow cooling rate. At approximately 910° C. or lower, γ-Fe transforms into α-Fe, which deteriorates pulverization efficiency during production of sintered magnets. If α-Fe remains even after sintering, magnetic characteristics of the sintered product are deteriorated. Thus, the ingot obtained through mold casting must be subjected to homogenization treatment at high temperature for a long period of time in order to remove α-Fe.
In order to solve the above problems; i.e., segregation of R-rich phase and precipitation of α-Fe, the strip casting method (abbreviated as SC method) and the centrifugal casting method (abbreviated as CC method), which ensure a cooling rate during casting of R—T—B alloy that is faster than that attainable by mold casting, are proposed and employed in actual production steps.
In the SC method, molten alloy is poured onto a rotatable copper roller for casting, the inside of which is cooled by water, and is formed into a strip having a thickness of about 0.1 to about 1 mm. During casting, the molten alloy is solidified through rapid cooling, to thereby prevent precipitation of α-Fe which is formed during mold casting and yield an alloy having a microcrystalline structure in which R-rich phase is minutely dispersed. Since the R-rich phase is minutely dispersed in the alloy ingot produced through the SC method, dispersion of R-rich phase in the product obtained by pulverizing and sintering the alloy also becomes satisfactory, to thereby successfully produce magnets of improved magnetic characteristics (Japanese Patent Application Laid-Open (kokai) Nos. 5-222488 and 5-295490).
Meanwhile, the CC method includes feeding a molten metal into the interior of a cylindrical mold which is rotating, to thereby simultaneously deposit and solidify the molten metal. Thus, the method attains an intermediate solidification rate between that attainable by mold casting and that attainable by the SC method (U.S. Pat. No. 2,817,624). This particular solidification condition is confirmed to be effective for producing a boundary phase alloy for employment in the two-alloy blending method (U.S. Pat. No. 3,234,741).
As compared with the mold casting method, the SC method and CC method attain high uniformity in microstructure. The uniformity in microstructure can be evaluated in terms of crystal grain size and dispersion state of R-rich phase as well as presence of precipitated α-Fe. When an alloy is cast through the mold casting method, a portion of the resultant alloy ingot which has remained in the vicinity of the mold and has been rapidly cooled exhibits a microstructure formed of minute equiaxed crystal grains called chill crystals and contains a comparatively finely dispersed R-rich phase. However, in the center portion of the alloy ingot where solidification is finally complete, crystal grains have a large grain size and R-rich phase forms aggregates in some regions, because of a considerably slow solidification rate in the center portion.
In the alloy flakes produced through the SC method, chill crystals may be formed on the side which has been in contact with a rotating roller for casting (hereinafter referred to as a mold side). However, an appropriately minute and uniform microstructure can be generally obtained through rapid-cooling solidification. In addition, since formation of α-Fe is suppressed, uniformity in R-rich phase contained in a sintered magnet, a final product, is enhanced, thereby preventing impairment of α-Fe in terms of pulverizability and magnetic characteristics.
When a molten alloy is cast through the CC method, the molten alloy is gradually deposited and the thus-solidified thin layers are stacked. Therefore, the cast product can possess a microstructure which is almost uniform from the mold side to the free surface side, despite its large thickness, except that chill crystals are formed in a mold side portion. However, since a conventional CC method (e.g., a method disclosed in U.S. Pat. No. 2,817,624) employs a comparatively high rate of feeding molten alloy, the substantial solidification rate becomes slower than that employed in the SC method. Thus, the conventional CC method attains an effect for preventing precipitation of α-Fe to a degree, which lies between that attainable by mold casting and that attainable by the SC method.
In recent years, Nd in R—T—B alloys for producing rare earth magnets has often been partially substituted by Pr. This is because partial substitution of Nd by Pr causes only a small variation in characteristics; Pr is less expensive than Nd; and production cost can be reduced. In the case of an R2Fe14B compound, saturation magnetization at room temperature of the compound (R═Nd) is known to be approximately 4% higher than that of the compound (R═Pr), but anisotropic magnetic field of the compound (R═Pr) is known to be approximately 5% higher than that of the compound (R═Nd). Regardless of whether R is Nd or Pr, phase conditions around R2Fe14B compound are substantially the same. Thus, even when Nd of R2Fe14B is partially substituted by Pr, phase constitutions remain substantially unchanged, and magnetism is not deteriorated by such a subtle change in microstructure.
The present invention is constituted by four aspects and respective aspects have following problems to be solved.
The problems to be solved by the first aspect of the present invention is described below.
From the viewpoint of cost reduction and effective utilization of resources, substitution in terms of R in R—T—B alloys for producing rare earth magnets, i.e., partial substitution of Nd by Pr, has been widely employed. However, the Pr content of R can be elevated up to about 10% by mass, because Pr is chemically active as compared with Nd. Such high chemical activity causes problematic oxidation during production of magnets or in the produced magnets.
As compared with the single-alloy method, the two-alloy blending method, which is widely employed for producing high-performance magnets, imposes more severe limitation on the amount of Pr added. The two-alloy blending method employs two types of raw material alloys; i.e., a main phase alloy, which predominantly provides R2Fe14B phase (main phase) and has a composition similar to that of R2Fe14B, and a boundary phase alloy, which predominantly provides R-rich phase (grain boundary phase) and has a TRE (Total Rare Earth content) greater than that of the main phase alloy.
In the two-alloy blending method, Pr is preferably added to the main phase alloy. When Pr is added to the boundary phase alloy containing a large amount of R-rich phase, which per se is prone to oxidation, activity is further increased. Thus, oxidation occurs predominantly during pulverization involved in magnet production steps and in the resultant micro-powder, leading to requirement of strong countermeasures for preventing oxidation, or deterioration in magnet characteristics caused by an increase in oxygen content thereof. Such countermeasures render the steps and apparatus for producing magnets complex, resulting in increased cost. In contrast, when Pr is added to the main phase alloy, Pr is predominantly incorporated into R2Fe14B phase, which per se is highly anti-corrosive, so that problematic oxidation can be mitigated. In addition, when Nd is partially substituted by Pr, an anisotropic magnetic field of R2Fe14B phase slightly increases. Thus, the micro-powder can be readily caused to be oriented during orientation in a magnetic field, thereby increasing magnetization and a degree of orientation of produced magnets.
As mentioned above, Pr is preferably added to the main phase alloy. However, in the course of partial substitution of Nd of the main phase alloy having a low TRE by Pr, α-Fe is prone to precipitate. One possible reason is that the substitution by Pr increases the difference between a temperature of the liquidus at which formation of γ-Fe (high-temperature phase) is initiated and the peritectic temperature at which formation of R2Fe14B phase is initiated. Since α-Fe is difficult to pulverize, efficiency of pulverization in magnet production steps is deteriorated, thereby reducing productivity of magnets. If unpulverized α-Fe remains in a pulverization apparatus, the composition of the resultant micro-powder varies. If α-Fe remains in a magnet even after sintering, magnetic characteristics of the magnet are considerably deteriorated.
According to the SC method, molten metal can be supercooled at high cooling rate to a temperature lower than peritectic temperature at which R2Fe14B phase is formed, thereby preventing precipitation of α-Fe. However, when an Nd—Fe—B ternary main phase alloy has an Nd content of about 28.5% by mass or less, sufficient supercooling cannot be attained, whereby α-Fe is formed. In addition, when Nd is partially substituted by Pr, precipitation of α-Fe is further promoted. Thus, in order to prevent precipitation of α-Fe, the TRE of the main phase alloy must be increased. In the two-alloy blending method, the TRE of the main phase alloy is preferably adjusted to as low a level as possible so as to enhance the mixing ratio of the boundary phase alloy.
An increase in B content is known to be remarkably effective for preventing precipitation of α-Fe. However, when the B content of the main phase alloy increases, the B content of the boundary phase alloy must be lowered in order to adjust the total B level of the finally produced magnet. Addition of Co or a heavy rare earth element to the main phase alloy is also effective for preventing precipitation of α-Fe. However, when the above compositional control approaches are employed, the degree of freedom in compositional design for magnet alloy decreases. Even when the two-alloy blending method is employed, an optimum combination of the compositions is difficult to attain.
The element Co, which exerts excellent effect for improving corrosion resistance, is preferably added to the boundary phase alloy (Kusunoki et al., T. IEEE Japan, Vol. 113-A, No. 12, 1993, 849-853). A heavy rare earth element is also confirmed to exert excellent effect for enhancing coercivity when the element is added to the boundary phase alloy (Ito et al., Journal of the Japan Institute of the Metals, Vol. 59, No. 1 (1995), 103-107).
The problems of the second aspect of the present invention are as follows.
A series of studies were carried out on the relationship between the microstructure of the cast R—T—B alloy ingot and the behavior upon hydrogen decrepitation or micro-pulverization, and has found that control of the dispersion state of R-rich phase is more critical, for providing a sintered magnet alloy powder of uniform particle size, than control of the crystal grain size of the alloy ingot. The inventor has also found that a region in which dispersion state of R-rich phase is excessively minute (fine R-rich phase region) formed on the mold side of the alloy ingot is a more critical factor for controlling the particle size of magnet powder than adverse effects of chill crystals, which are in fact contained in an alloy ingot in amounts of some % or less. In other words, the inventor has confirmed that the percent volume of fine R-rich phase region may be in excess of 50% even when the number of chill crystals contained in the R—T—B alloy ingot is decreased through modification of the composition of the alloy ingot or production conditions; that the fine R-rich phase region deforms the particle size distribution of the magnet alloy powder; and that the fine R-rich phase region must be reduced in order to enhance magnet characteristics.
The problems to be solved by third aspect is described below.
Through the method disclosed in Japanese Patent Application No. 2001-383989, reduction of fine R-rich phase region and yielding of uniform microstructure can be attained to some extent. However, other than surface conditions of a roller for casting, there are a variety of factors which determine the microstructure, and such factors are difficult to completely control during actual R—T—B alloy production. Thus, fine R-rich phase region may be formed at a portion of the alloy.
The problems to be solved by the fourth aspect are as follows.
The present inventor has carried out extensive studies on the relationship between the microstructure of the cast R—T—B alloy ingot and the behavior upon hydrogen decrepitation or micro-pulverization, and has found that control of the dispersion state of R-rich phase is more critical, for providing a sintered magnet alloy powder of uniform particle size, than control of the crystal grain size of the alloy ingot. The inventor has also found that a region in which dispersion state of R-rich phase is excessively minute (fine R-rich phase region) formed on the mold side of the alloy ingot is a more critical factor for controlling the particle size of magnet powder than adverse effects of chill crystals, which are in fact contained in an alloy ingot in amounts of some % or less. In other words, the inventor has confirmed that the percent volume of fine R-rich phase region may be in excess of 50% even when the number of chill crystals contained in the R—T—B alloy ingot is decreased through modification of the composition of the alloy ingot or production conditions; that the fine R-rich phase region deforms the particle size distribution of the magnet alloy powder; and that the fine R-rich phase region must be reduced in order to enhance magnet characteristics.
The present invention has been accomplished on the basis of this finding and an object thereof is to provide a method for producing an rare-earth-containing alloy flake, the method more effectively preventing formation of fine R-rich phase region in a cast rare-earth-containing alloy ingot made of an R—T—B alloy, and a rare-earth-containing alloy flake having a structure with excellent uniformity produced by the above method.